Assessment of CuO thin films for its suitablity as window absorbing layer in solar cell fabrications

Materials Research Bulletin 68 (2015) 1–8

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Assessment of CuO thin films for its suitablity as window absorbing layer in solar cell fabrications R. Shabu a , A. Moses Ezhil Raj a, * , C. Sanjeeviraja b , C. Ravidhas c a b c

Department of Physics and Research Centre, Scott Christian College (Autonomous), Nagercoil 629 003, India Department of Physics, Alagappa Chettiar College of Engineering & Technology, Karaikudi 630 004, India Department of Physics, Bishop Heber College (Autonomous), Tiruchirapalli 620 017, India

A R T I C L E I N F O

A B S T R A C T

Article history: Received 18 August 2014 Received in revised form 12 February 2015 Accepted 5 March 2015 Available online 14 March 2015

As major attention has been paid to solar cells with extremely thin absorber, present study embark on the structural, surface, optical and electrical characterization of spray deposited CuO absorber layers prepared using precursors of different concentrations at the constant substrate temperature of 350
C. Confirmed monoclinic lattice and the Cu—O vibrational modes implied the tenorite phase formation of CuO. Smooth surface and the uniform distribution of a surface grain which contains Cu2+ as the main oxidation state of the cations established the suitability of the CuO thin films as energy absorbers. Optical absorption in selective regions, estimated optical bandgap and low intense photoluminescence peaks favored respectively the easy charge transfer and the light retaining capacity of the CuO films. Probably, a combination of electrical conduction with nanostructured CuO will be helpful for future developments of substrates for solar cells with CuO as absorber layers. ã 2015 Elsevier Ltd. All rights reserved.

Keywords: A. Thin films B. Chemical synthesis C. X-ray diffraction C. Raman spectroscopy D. Electrical properties

1. Introduction Solar selective absorber would be a material that can absorb solar radiation (a = 1) in the visible and near infrared region (0.3–2 mm) without emission (e = 0) in the infrared region (2–20 mm) of the solar spectrum in order to fully utilize the high energy radiation as well as to minimize undesired thermal losses [1–4]. Such an ideal material is not attainable, but it is possible to produce selective surfaces having high solar absorptance and low emittance for long wavelengths [5,6]. In addition, long term stability and desired spectral selectivity are also the important requirements for environmentally friendly absorber coatings. Even though it is a challenging task to develop solar energy converting devices using low cost techniques, several solar selective coatings have been developed using semiconductor oxides [7]. Among the various semiconducting oxides under research, copper oxide is attractive as a selective solar absorber since it has high solar absorbency and a low thermal emittance [8–10]. Copper oxide has two phases: cuprite (Cu2O) and tenorite (CuO). It has been reported that Cu2O has a cubic structure with a bandgap of about 2.0–2.6 eV [11,12], while CuO has a monoclinic structure and a bandgap of 1.3–2.1 eV [12]. The p-type conduction property of

* Corresponding author. Tel.: +91 4652 232888; fax: +91 4652 229800. E-mail address: [email protected] (A. Moses Ezhil Raj). http://dx.doi.org/10.1016/j.materresbull.2015.03.016 0025-5408/ ã 2015 Elsevier Ltd. All rights reserved.

copper oxide is mainly attributed to the negatively charged Cu vacancies [13–16]. Amongst the mono oxides of 3d transition series elements, CuO is unique as it has a square planar coordination of copper by oxygen in the monoclinic structure (s.g.: C62h ðC2=cÞ). Copper oxide thin films have wide range of applications in energy harvesting, storage and solar selective absorbers. The use of single layer of black copper/nickel films based on copper oxides and nickel oxides for solar coatings was reported earlier [17]. Recent research efforts on CuO based solar cell in the configuration of p-CuO/n-ZnO:Sn was developed and its energy conversion efficiency was reported (h = 0.232%) [18]. Fabrication of the solar cell with configuration (ITO/CuO/ZnO/Al) was already reported by H. Kidowaki et al. [19]. The selective coatings evolved from sputter and rapid thermal annealing techniques for the formation of p-CuO/n-Si heterojunction solar cells have achieved an overall power conversion efficiency of 0.36% [20]. These advanced coatings require complex and high cost equipments and provide a rather poor efficiency. Moreover in most of the techniques, surfaces with mixed phases like Cu, CuO and Cu2O, are generally obtained at room temperature. Conversion efficiency mainly depends on surface morphology and roughness that plays an important role in improving the absorber efficiency. Spray pyrolysis deposition provides textured/ rough surfaces that can favor multiple internal reflections leading to increased values of solar absorption [21]. Therefore chemical spray pyrolysis deposition has been engaged in the present work to

(c)

(b)

(1 1 1)

prepare phase pure nanostructured CuO thin films on glass substrates. The paper focuses the properties of the films and the results are correlated for its suitability as an absorbing layer in photovoltaic devices.

(-1 1 1)

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(-1 1 1)

2

0.2 M

3. Results and Discussion 3.1. X-ray diffraction studies Fig. 1 shows the XRD pattern of prepared CuO thin films at a constant substrate temperature of 350
C for three different precursor concentrations (0.10 M, 0.15 M and 0.20 M). Prepared thin film sample using 0.1 M solution is amorphous, whereas films prepared using higher precursor concentrations are crystalline with systematic peaks at specified 2u locations. Variations in the crystallinity may be due to the thickness of the films. Thickness of the CuO film deposited using 0.1 M precursor is 1.03 mm, which has less number of crystalline planes to have diffraction peaks and are therefore amorphous. However, the thickness of the CuO films deposited using 0.15 and 0.2 M precursors are respectively 1.54 and 1.94 mm having more lattice planes and exhibit intense diffraction peaks. Two prominent peaks were observed at 35.50
and 38.73


(1 1 1)

CuO thin films were prepared by the chemical spray pyrolysis (CSP) technique by spraying 0.1 M copper acetate monohydrate [Cu(CH3COO)2H2O (Sigma–Aldrich)] on to preheated substrates placed inside a tubular furnace at 350
C. The solution flow rate was controlled by a flow meter (5 ml min1) and compressed purified air was used as the carrier gas (0.4 kg/cm2). Distance between the nozzle and the substrate was maintained at 30 cm for all the depositions. Coatings were repeated for two more precursor concentrations (0.15 M and 0.20 M) and the thickness of the prepared film was measured using Mitutoya Surftest SJ-301 stylus type surface profiler. Microstructure and crystallinity of the films were characterized using PANalytical-3040 X’pert Pro X-ray diffractometer with copper target (l = 1.5405 Å) in u–2u scan mode. Continuous scanning was applied with a slow scanning speed (step size = 0.01
) and small time constant (step time = 3 s). Raman spectra were recorded in the spectral range 100–1500 cm1 using an FT-Raman spectrometer (Bruker RFS 27-Multi RAM, Germany) with the excitation of laser source Nd:YAG having the excitation wavelength of 1064 nm. Surface analysis was performed through X-ray photoelectron spectroscopy (XPS) using Kratos Analytical, Ultra axis machine (calibrated) under standard protocols (Monochromatic Al Ka X-ray hv = 1486.6 eV, incident angle of 30
with respect to surface normal, collection of photoelectrons at a takeoff angle of 50
with respect to surface normal, analysis area of 400 nm in diameter and analysis depth of 10 nm). The carbon correction was done using manufacturer’s standard software. The surface morphology of the films was investigated by AFM (Park systems: XE 100) in the non-contact mode at a 285 kHz resonance frequency and an approximate 42 N/m constant force. AFM micrographs were recorded from different regions of the samples for the sampling areas of 5 mm  5 mm. Surface morphology of films was also observed through scanning electron microscopy (JEOL Model JSM - 6390LV). Optical properties of the samples were investigated using a UV–vis–NIR double beam spectrophotometer (Varian Carry 5000). Photoluminescence spectra were recorded with a luminescence spectrometer (Perkin-Elmer, LS 45) having a spectral resolution of 0.5 nm. Standard two probe technique was used for the measurement of electrical conductivity using precision electrometers.

Intensity (a.u.)

2. Experimental details 0.15M

(a) 0.1 M

10

20

30

40

50

60

70

80

2θ (degrees) Fig. 1. XRD pattern of prepared CuO thin films for three different precursor concentrations (a) 0.10 M, (b) 0.15 M and (c) 0.20 M.

for the films using 0.15 M and 0.20 M solution, which corresponds to the ð111Þ and (111) diffraction planes with interplanar distances 2.528 Å and 2.324 Å, respectively. As the concentration was increased to 0.2 M, the intensity of the prominent peaks also increased which is consistent with the results reported by Gopalakrishna et al. [27]. Obtained values are in good agreement with standard data of CuO phase (JCPDS Card No.: 80-1916). In comparison to this, Balamurugan and Mehta [11] reported a mixture of Cu2O and CuO phase by an activated reactive evaporation technique. Nair et al. [22] could only achieve a deposition of Cu phase along with a minority Cu2O phase at room temperature by the chemical deposition technique. However in the present study, spray pyrolysis technique with specially designed spray nozzle and the optimized deposition parameters yielded only the single phase tenorite CuO phase with monoclinic structure at the lowest deposition temperature of 350
C. Obtained observations are reliable with the results reported by Gopalakrishna et al. [27], Morales et al. [24], Senthil Kumar et al. [25] and Cho et al. [26] for CuO films prepared by spray pyrolysis technique. In both the films, ð111Þ is the predominant crystallographic plane indicating a perpendicular alignment of the c-axis of the grains [23]. The mean crystallite size (D) of nano crystalline CuO thin films was calculated using Scherrer's formula (Eq. (1)), neglecting the peak broadening due to residual stresses in the films [35]. D¼

0:94l b cosu

(1)

where, ‘l’ is the wavelength of the X-rays; ‘u ’ is the Bragg’s diffraction angle, and ‘b’ is the full width at half maximum (FWHM) of the diffraction peaks (in radians). The strain (e) developed in films due to lattice misfit, and the dislocation density (d) defined as the length of dislocation lines per unit volume of the crystal, number of crystallites was evaluated using Eqs. (2) and (3) respectively [28].

e¼

d¼

b cosu 4

1 D2

(2)

(3)

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Table 1 Structural parameters of the CuO thin films prepared using precursor of different concentrations. CuO thin films prepared using precursor concentration

Microstructural parameters

a(Å) b(Å) c(Å) b(
) Cell volume (Å)3 Micro strain (e)  103 Crystallite size (nm) Dislocation density (1015 lines/m)

0.15 M

0.20 M

4.692 3.429 5.137 99.5 81.51 0.973 37 0.730

4.685 3.433 5.149 99.6 81.66 0.695 52 0.369

The lattice parameters (a 6¼ b 6¼ c,a = g = 90
6¼ b for monoclinic structure) and the unit cell volume were calculated using Eqs. (4) and (5), ! 2 2 2 1 1 h k sin b 12 2 h l cosb ¼ þ þ  (4) 2 2 ac c2 sin2 b a2 b d

group symmetry of C2h 6 [30,31]. The metal cation copper forms four coplanar bonds with oxygen, which again self coordinated by four copper atoms, forming two sets of chains directed along (11 0) direction [31–40]. The factor group analysis yields [32],

V ¼ abcsinb

O: G = Ag + 2Bg + Au + 2Bu

(5)

Calculated lattice parameter and other related structural informations are listed in Table 1. Obtained lattice parameters are comparable to the reported results [a = 4.6927(4), b = 3.4283(4), c = 5.137(6) and b = 99.546(9)]. Moreover, the crystallite size increases with the increase in film thickness and as a result, both the strain and dislocation density decreases. The decrease in dislocation density and strain indicates the decrease in the concentration of lattice imperfections, and better crystallization of high quality CuO thin films suitable for photovoltaic applications. The preferential orientation of the deposited film was estimated from the texture coefficient (TC): TCðhklÞ ¼

 1 I ðh k lÞ 1 n I ðh k lÞ S I0 ðh k lÞ n i¼1 I0 ðh k lÞ

(6)

where, ‘I0’ represents the intensity of reference diffraction pattern (JCPDS card no. 80-1916), ‘I’ is the observed intensity of the (h k l) plane and ‘n’ is the reflection number. Usually the texture coefficient value is one for the random distribution of grains in all the reflecting planes of the film, whereas its values greater than one indicate the abundance of the grains in a given (h k l) plane. For the TC(h k l) values that lie between zero and one indicate the lack of grain orientation in that direction [29]. Calculated texture coefficient for the CuO films is listed in Table 2. It is evident that the

Cu: G = 3 Au + 3Bu

Gtot CuO ¼ 4Au þ 5Bu þ Ag þ 2Bg Of these, 1Au and 2Bu are the acoustical modes, so that the total vibrational modes and their activity are [33–35]:

Gvib CuO ¼ 3Au þ 3Bu þ Ag þ 2Bg Thus, three modes are Raman active (Ag, Bg) and six infrared (Au, Bu) active modes are to be expected in the spectra of CuO [36,37]. Fig. 2 shows the FT- Raman spectra of CuO thin films coated on glass substrate using two precursors of different concentrations. As oxygen atoms are involved in the Raman active modes, Raman shift of Ag modes are less compared to Bg, as is in accordance with experimental results 152.64 cm1 (A1g) and 384.2 cm1 (B1g). Another peak lies at 570.69 cm1 corresponds to the B2g mode of CuO. Observed locations of the peaks are comparable to the previously reported results [38–42]. The band at about 1093.71 cm1 can be assigned to multi-phonon transition [30,31].

grains are abundantly oriented along the ð111Þ plane and also the TC value of this particular plane increases with film thickness. 3.2. FT-Raman spectral studies CuO crystallizes in a monoclinic lattice with four formula units in the crystallographic primitive cell and belongs to the space Table 2 Calculated texture coefficient values of CuO thin films. Precursor concentration

2u (
)

d (Å)

hkl

I (h k l)

I0 (h k l)

TC (h k l)

0.10 M

–

–

–

–

–

–

ð1 1 1Þ (1 1 1)

100.00

95.4

1.16

75.28

99.9

0.83

ð1 1 1Þ (1 1 1)

100.00

95.4

1.25

93.04

99.9

1.11

0.15 M

0.20 M

35.50

2.528

38.73

2.324

35.40

2.535

38.59

2.332

Fig. 2. Raman spectra of CuO thin films on glass substrate at two different concentrations.

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3.3. Surface characterization 3.3.1. XPS analysis The XPS survey spectrum of the CuO thin films prepared using 0.2 M copper acetate is presented in Fig. 3. The peaks corresponding to Cu 3p, Cu 3s, Cu 2p and Cu Auger, O 1s and O Auger are obviously observed. The high-resolution spectra corresponding to Cu 2p and O 1s are given in Fig. 3a and b. They are plotted after correction of charging effects using the binding energy of 284.8 eV corresponding to the C 1s peak. Fig. 3a shows the high resolution Cu 2p3/2 and 2p1/2 spectra. The main broader peak (3.27 eV) of Cu 2p3/2 lies at 933.38 eV, which can be assigned to Cu2+ ions, rather than Cu+ ions, which have a main peak at a lower binding energy of about 932.4 eV [43]. On peak fitting the Cu 2p3/2 peak, absence of peaks below the binding energy 933.38 eV clearly indicates the formation of pure CuO nanostructures without the Cu2O phase. In addition to the main peaks of Cu 2p3/2, their satellite peaks are observed on the high-binding-energy side more than 9 eV. This satellite is a characteristic of materials having a d9 configuration in the ground state, such as copper dihalides [44] or CuO [43]. The structure seen in the satellite line is due to the multiplet splitting in the 2p5 3d9 final state [45]. Usually the spectrum of Cu2O has only one peak at 932.4 eV, which is significantly narrower (1.9 eV). Also for Cu2O with a completely filled shell (d10), the satellite peak is absent because screening via a charge transfer into the d states is not allowed. Moreover, in the sample, the main peak of 2p1/2 is located at a binding energy that is 20 eV higher than that of their 2p3/2 main peak which is the characteristic for the formation of CuO. Based on the above-described results, Cu2+ is the dominant oxidation state at the film surface.

In the O 1s spectrum (Fig. 3b) after curve deconvolution, there are three components corresponding to various chemical states. The main peak is located at 531.84 eV and the FWHM of the peak is about 2.3 eV which is large than that of single crystal cupric oxide (0.8 eV), where the satellite components almost disappears [46]. The energy of one of the satellites in the sample is centered at approximately 4.4 eV higher than their main peak. It is well-known that the main peak and satellites of O 1s spectra correspond to different chemical states of the oxygen atoms [47]. The main peak at lower binding energy is attributed to the normal O2, interacting with the copper atoms to form the chemical Cu—O bond. The satellite at higher energy is ascribed to the extra lattice oxygen, indicating non-stoichiometric nature of the films [46]. 3.3.2. AFM analysis Fig. 4 shows the typical three dimensional (3D) and two dimensional AFM images of the CuO films deposited on glass substrates using two different precursor concentrations 0.15 M and 0.20 M. Images revealed the “hills and valley” like surface evolution of the films which are uniformly distributed over the entire substrate surface. Also the grains are vertically aligned and its approximate sizes are respectively 160 nm and 200 nm for the two films. From the 3D images it can be understood that the height of the hills increases and width of the valley decreases when the precursor concentration increases. So that the root mean square roughness decreases with increase of precursor concentration and correspondingly film quality increases. This may be due to the agglomeration of nano sized crystallites when the thickness of the film increases. The surface roughness of the films were estimated from the histogram drawn using the software WSxM, and the root mean square value decreases to 21 nm for films coated using

Fig. 3. XPS survey spectrum for the CuO thin film prepared using 0.2 M precursor. (a) Cu 2p XPS core level spectrum showing the surface dominated by CuO. (b) Deconvoluted profile for O 1s core level spectrum of CuO thin film sample.

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Fig. 4. 2D and 3D images of CuO thin film prepared using precursor of different concentrations: (a) 0.15 M and (b) 0.20 M.

0.20 M solution, whereas it is 43 nm for the films prepared using 0.15 M solution. These types of films having vertically aligned grains favor multiple internal reflections that lead to more optical absorption. 3.3.3. SEM analysis The SEM images show the grain formation, significantly depending on the precursor concentrations (Fig. 5). It can be seen that the prepared thin films have uniform surface morphology over the entire substrate and are of good quality. In Fig. 5a, the surface is with holes which indicate the commencement of particle formation and therefore films are amorphous in nature as revealed by XRD studies. Fig. 5b and c shows that the particles begin to grow in size and dense with increase of precursor concentration and therefore the crystallinity of the films improves gradually. The film deposited at 0.2 M concentration is uniform, dense and well adhered to the glass substrate. This kind of dense and compact morphology is essential for the applications in thin film solar cells, to prevent the leakage of photo-current [48]. 3.4. Optical properties 3.4.1. Optical absorption and transmission. Fig. 6 shows the optical absorption spectra of nanocrystalline CuO thin films prepared for two different thicknesses using

precursors of different concentrations. The films exhibit high absorption in the solar spectrum region (0.2–0.8 mm), the reflection is very low and the reflectance in some region is near to 0%, while in the infrared region (0.8–2.5 mm), the reflectance is rather high. This indicates that the CuO thin films have good solar selectivity. The transmission spectra were recorded at room temperature in air to obtain information on the optical properties of the cupric oxide thin films (Fig. 7). It is observed that the films show different transmissions for various wavelengths. The overall transmittance of the film is very low and this may be due to the observed solar selectivity of these thin films are greater. The optical parameters such as absorption coefficient and bandgap are determined from optical absorption measurements. The value of absorption coefficient for strong absorption region of thin film is calculated using the equation [49]:

a¼

Aðhv  Eg Þn hv

(7)

where, ‘Eg’ is the bandgap energy, ‘A’ is a constant, and ‘n’ is equal to ‘1/2’ and ‘2’ for allowed direct and indirect transitions respectively. The studies indicate direct bandgap, and this property is suitable for photovoltaic applications. The absorption region threshold which is defined by the bandgap of the materials can be estimated from the Tauc plot (Fig. 6). Estimated bandgap of CuO thin films

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Fig. 5. SEM micrograph of CuO thin film deposited using precursors of different concentrations: (a) 0.10 M, (b) 0.15 M and (c) 0.20 M.

was found to decrease from 2.55 eV to 2.50 eV with the increase in precursor concentration. Compared with bulk CuO, the bandgap of nanostructured CuO is blue shifted, with reported values ranging widely from 1.2 eV to 2.1 eV [50]. Other researchers have also reported larger bandgap up to 4.13 eV for 10 nm quantum dots [51] and 3.02 eV for well-aligned arrays of CuO nano platelets [52]. These variations in bandgap are due to the quantum confinement effect. The reduction in particle sizes results in increase of surface/volume ratio. As a result surface atom has lower coordination number and atomic interaction which increases the highest valance band energy and decreases the lowest unoccupied conduction band energy. This leads to increase

in bandgap energies [53]. Therefore, CuO absorbs strongly throughout the visible spectrum with a slight transparency for bigger bandgap nanostructured samples, which absorb in the UV region and reflect in the IR region.

Fig. 6. Optical absorption spectra of CuO thin films prepared using precursors of different concentrations. Inset shows reflectance spectra.

Fig. 7. Optical transmittance spectra of CuO thin films and the inset shows the constructed Tauc plot for the estimation of optical bandgap energy.

3.4.2. Photoluminescence Photoluminescence (PL) spectra of CuO thin films prepared using two different precursor concentrations (0.15 M and 0.20 M) are illustrated in Fig. 8. During excitation with wavelength of 375 nm, room temperature PL spectra of the samples revealed similar features with five major peaks (482, 492, 509, 613 and 657 nm) correspond to various emissions of CuO. Since the band

R. Shabu et al. / Materials Research Bulletin 68 (2015) 1–8

613

1000000

PL Intensity (a.u.)

7

0.15M 0.20M

800000 657

509 600000 400000

482

492

200000 0 450

500

550

600

650

700

Wavelength (nm) Fig. 8. PL spectra of CuO thin films prepared using precursors of different concentrations.

structure calculations of CuO nanostructures are limited even in its bulk form due to large variations in the electron and hole effective mass, prediction and PL band assignments are complex [51,54–56]. However, PL bands in the blue region (482 and 492 nm) are caused by transition vacancy of oxygen and interstitial oxygen [57]. The PL peak at 509 nm, which corresponds to green emission, arises from the singly ionized oxygen vacancy [57]. Existence of oxygen vacancies in the prepared samples may be due to the incomplete oxidation taking place at relatively low temperature deposition at 350
C [58]. It could be the reason for the intense emission of green band. PL bands at the red region (613 and 657 nm) are the deep level emissions related to the interstitials in CuO [59]. Moreover, on increasing the precursor concentration the peak intensity drastically decreases that indicates the defect free nature of the prepared nanostructures. PL bands corresponding to Cu2O are not observable in both the cases as confirmed through XRD results. Low intense emission peaks again reveal the light retaining capacity of the deposited films, which is an imperative character required for solar cell window layers. 3.5. Electrical properties. The temperature dependent electrical conductivity (s ) of the CuO thin films for two different thicknesses (1.04 and 1.54 mm) were measured in the range 300–473 K. The room temperature electrical conductivity of the film is respectively 6.37  105 and 3.89  104 mho/m and the conductivity increases with thickness. The observed lesser conductivity in thinner films can be explained by the lower degree of crystallinity and the smaller crystallite size, as detected by the XRD analysis. The influence of crystallite size on activation energy has been studied extensively [60]. It perceives that the conductivity also increases with temperature indicating the semiconducting behavior of the films. The activation energy of the deposited films has been calculated from the local gradient of ln s versus 1000/T plots, based on the following equation:   E (8) s ¼ s 0 exp  a kT where, Ea is the activation energy of conduction, s 0 is a temperature independent factor, k is the Boltzmann’s constant and T is the absolute temperature. The electrical conductivity (s ) as a function of the reciprocal temperature is shown in Fig. 9. The electrical results show temperature dependence conductivity and exhibit single type of conduction channel in the measured temperature range. Usually, below room temperature the electrical conductivity (s ) is insensitive to temperature and corresponds to transport by variable range hopping conduction

Fig. 9. Plot of ln s versus 1000/T for CuO thin films prepared for two thicknesses using precursors of different concentrations.

(VRH) in the localized states near the Fermi level. In polycrystalline materials, the VRH conduction process exists in the grain boundaries where the carriers do not have sufficient energy to cross the potential barrier and transfer themselves into the grain by the process of thermionic emission. On the other hand, the exponential increase of (s ) with T can be regarded as regular band-type conduction in extended state. The activation energy for conduction (Ea) is respectively 1.14 and 0.97 eV for the CuO films prepared using 0.15 M and 0.20 M precursors. From the observed results, it was found that the activation energy is thickness dependent and decreases with increasing the thickness. This can be explained by the polycrystalline nature of the films, which is detailed with a model established by Seto [61]. A polycrystalline film material contains a large number of micro crystallites separated by grain boundaries that plays an important role between crystallites in determining the conductivity of a polycrystalline film. The incomplete atomic bonding at a grain boundary can act as trap centers. These centers trap the charge carriers there, to build a local space charge. This charge then impedes the transition of charge carriers from one crystallite to another. Thus, the values of Ea will vary with applied voltage, which is about different trapping levels situated between the valence and conduction bands. As the thickness decreases, the crystallite size decreases and this leads to an increment in the trapping states at grain boundary. Trapping states are capable of trapping free carriers and, as a consequence, more free carriers become immobilized as trapping states increase. In other words, larger crystallite size results in a lower density of grain boundaries, which behave as traps for free carriers and as barriers for carrier transport in the film. Hence, an increase in grain size can cause a decrease in grain boundary scattering, which leads to an increase in conductivity. This is because of improved crystallinity and the increase grain size of the films as the film thickness increases [62,63]. 4. Conclusion CuO thin films have been successfully deposited on commercial flat glass substrates at 350
C by the chemical spray pyrolysis technique using precursors of different concentrations. Precursor concentration produced significant variations in film thickness and the XRD patterns showed an improvement in crystallinity with increasing film thickness. Prepared CuO films were polycrystalline

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and monoclinic in structure without the additional phases like Cu2O. The mono phase formation was confirmed through the FT-Raman spectra and photoelectron spectroscopy. The Cu 2p1/2 was located at a binding energy higher than 20 eV that of their 2p3/2 main peak, confirmed the characteristic formation of CuO. Even though the surfaces of films were smooth, poor transmittance revealed its high absorptance and reflectance. Emission bands in PL spectra were defect related and these defects were reduced much on improving the thickness of the films. As the electrical conductivity and activation energy were observed to be thickness dependent, the thickness effect was associated with the change in grain boundary potential with grain size, according to Seto’s model. Acknowledgement One of the authors (R. Shabu) is grateful to CSIR, Government of India, for the award of Senior Research Fellowship under CSIR–Emeritus Scientist Scheme. References [1] J.A. Duffie, W.A. Beckman, Solar Engineering of Thermal Processes, John Wiley, New York, 1980. [2] K.W. Boer, Advances in Solar Energy, Plenum Press, New York, London, 1990. [3] K.D. Lee, W.C. Jung, J.H. Kim, Sol. Energy Mater. Sol. Cells 63 (2000) 125–137. [4] N.C. Bhowmik, J. Rahman, M.A.A. Khan, Z.H. Mazumder, Renewable Energy 24 (2001) 663–666. [5] P.S. Patel, L.D. Kadam, C.D. Lokhande, Thin Solid Films 272 (1996) 29–32. [6] F. Jahan, M.H. Islam, B.E. Smith, Sol. Energy Mater. Sol. Cells 37 (1995) 283–293. [7] R. Motoyoshi, T. Oku, A. Suzuki, K. Kikuchi, S. Kikuchi, B. Jeyadevan, J. Cuya, Adv. Mater. Sci. Eng. 2010 (2010) 562842-1-11. _ Erdog an, Ö. Güllü, J. Alloys Compd. 492 (2010) 378–383. [8] Y. I. [9] T. Maruyama, Sol. Energy Mater. Sol. Cells 56 (1998) 85–92. [10] A. Chen, G. Yang, H. Long, F. Li, Y. Li, P. Lu, Thin Solid Films 517 (2009) 4277–4280. [11] B. Balamurugan, B.R. Mehta, Thin Solid Films 396 (2001) 90–96. [12] G. Papadimitropoulos, N. Vourdas, V.Em. Vamvakas, D. Davazoglou, Thin Solid Films 515 (2006) 2428–2432. [13] H. Raebiger, S. Lany, A. Zunger, Phys. Rev. B 76 (2007) 045209 (1–5). [14] O. Porat, I. Riess, Solid State Ionics 81 (1995) 29–41. [15] A.O. Musa, T. Akomolafe, M.J. Carter, Sol. Energy Mater. Sol. Cells 51 (1998) 305–316. [16] A.A. Ogwu, E. Bouguerel, O. Ademosu, S. Moh, E. Crossan, F. Placido, J. Phys. D: Appl. Phys. 38 (2005) 266–271. [17] M. Voinea, E. Ienei, C. Bogatu, G.C. Chitanu, A. Duta, J. Nanosci. Nanotechnol. 9 (2009) 4279–4284. [18] E.O. Omayio, P.M. Karimi, W.K. Njoroge, F.K. Mugwanga, Int. J. Thin Films Sci. Technol. 2 (1) (2013) 25–28. [19] H. Kidowaki, T. Oku, T. Akiyama, J. Phys. Conf. Ser. 352 (2012) 012022. [20] S. Masudy-Panah, G.K. Dalapati, K. Radhakrishnan, A. Kumar, H.R. Tan, E.N. Kumar, C. Vijila, C.C. Tan, D.Z. Chi, Prog. Photovoltaics Res. Appl. (2014) , doi:http://dx.doi.org/10.1002/pip.2483. [21] I. Oja Acik, A. Junolainen, V. Mikli, M. Danilson, M. Krunks, Appl. Surf. Sci. 256 (2009) 1391–1394. [22] M.T.S. Nair, L. Guerrero, O.L. Arenas, P.K. Nair, Appl. Surf. Sci. 150 (1999) 143–151. [23] V. Dhanasekaran, T. Mahalingam, Mater. Res. Bull. 48 (2013) 3585–3593.

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